Nickel base super alloys and methods of making the same

ABSTRACT

Methods of forming an intermediate alloy and a Ni-base super alloy are disclosed along with the intermediate alloy and the Ni-base super alloy formed by the method. The method includes at least partially melting and solidifying a powder including about 5 to 15 wt. % of Co, 10 to 20 wt. % of Cr, 3 to 6 wt. % of Mo, 3 to 6 wt. % of W, 2 to 4 wt. % of Al, 4.2 to 4.7 wt. % of Ti, 0.01 to 0.05 wt. % of Zr, 0.015 to 0.060 wt. % of C, 0.001 to 0.030 wt. % of B and balance substantially Ni to form an intermediate alloy including a dendrite structure that includes columnar regions and intercolumnar regions and a primary dendrite arm spacing less than about 3 micrometers. The intermediate alloy is heat-treated to form the texture-free Ni-base super alloy.

CROSS REFERENCE TO RELATED APPLICATIONS

This patent application is a divisional of U.S. Non-Provisional patentapplication Ser. No. 14/963,366 filed on Dec. 09, 2015, now U.S. Pat.No. 10,378,087, which is incorporated by reference herein in itsentirety.

BACKGROUND

The present invention relates to nickel base (Ni-base) super alloys andmethods of making the same. More particularly, this invention relates tointermediate alloys and texture-free Ni-base super alloys having aparticular composition, and methods of making the same.

Ni-base super alloys are very useful set of alloys that can be designedto be used with substantial creep and oxidation resistances at hightemperatures, often in excess of 0.7 of their absolute meltingtemperatures. One form of high-temperature nickel base alloy ofparticular interest is a cast form, designed for desired creep andenvironmental properties, such as for example, oxidation resistance andhot corrosion resistance. One of the nickel base composition that iswidely used, especially in hot gas path components, is the Rene 80™material trademarked by General Electric Company.

Additive manufacturing is a suite of emerging technologies thatfabricate three-dimensional objects directly from digital models throughan additive process, typically by depositing layer upon layer andjoining them in place. Unlike traditional manufacturing processesinvolving subtraction (e.g., cutting and shearing) and forming (e.g.,stamping, bending, and molding), additive manufacturing joins materialstogether to build products.

Certain components, such as for example, turbine engine hot gas pathcomponents made by nickel base super alloys can benefit from the designflexibility of additive manufacturing techniques. However, additivemanufacturing methods may create certain challenges in themanufacturability of the nickel base super alloys and further in formingthe nickel base super alloys having properties similar to the nickelbase super alloys that are manufactured by the traditional castingmethod. Therefore, it is desirable to design a nickel base super alloythat can be manufactured using the additive manufacturing methods andhave properties that are similar to their cast counterparts.

BRIEF DESCRIPTION

Briefly, one embodiment is directed to a method. The method includes atleast partially melting and solidifying a powder including about 5 to 15wt. % of Co, 10 to 20 wt. % of Cr, 3 to 6 wt. % of Mo, 3 to 6 wt. % ofW, 2 to 4 wt. % of Al, 4.2 to 4.7 wt. % of Ti, 0.01 to 0.05 wt. % of Zr,0.015 to 0.060 wt. % of C, 0.001 to 0.030 wt. % of B and balancesubstantially Ni to form an intermediate alloy including a dendritestructure that includes columnar regions and intercolumnar regions. Aprimary dendrite arm spacing of the dendrite structure is less thanabout 3 micrometers. The method further includes heat-treating theintermediate alloy in a temperature range from about 1050° C. to about1250° C. to form a texture-free super alloy.

Another embodiment is directed to an intermediate alloy. Theintermediate alloy includes a dendrite structure that includes columnarregions and intercolumnar regions. A primary dendrite arm spacing of thedendrite structure is less than about 3 micrometers. The intermediatealloy has a composition that includes about 5 to 15 wt. % of Co, 10 to20 wt. % of Cr, 3 to 6 wt. % of Mo, 3 to 6 wt. % of W, 2 to 4 wt. % ofAl, 4.2 to 4.7 wt. % of Ti, 0.01 to 0.05 wt. % of Zr, 0.015 to 0.060 wt.% of C, 0.001 to 0.030 wt. % of B and balance substantially Ni.

Another embodiment is directed to a Ni-base super alloy. The Ni-basesuper alloy includes a composition that include about 5 to 15 wt. % ofCo, 10 to 20 wt. % of Cr, 3 to 6 wt. % of Mo, 3 to 6 wt. % of W, 2 to 4wt. % of Al, 4.2 to 4.7 wt. % of Ti, 0.01 to 0.05 wt. % of Zr, 0.015 to0.060 wt. % of C, 0.001 to 0.030 wt. % of B and balance substantiallyNi. The Ni-base super alloy includes a gamma phase matrix, precipitatesof a gamma-prime phase, and metal carbides precipitated from the gammaphase matrix. The metal carbides that are in the Ni-base super alloy isless than about 0.3 mole % of the composition, and an average size ofthe metal carbides is less than about 1 micrometer.

DRAWING

These and other features, aspects, and advantages of the presentinvention will become better understood when the following detaileddescription is read with reference to the accompanying drawing, wherein:

FIG. 1 illustrates a low magnification prior art microstructure image ofan as-cast alloy formed by using a baseline Rene 80 composition;

FIG. 2 illustrates a low magnification microstructure image of anas-built alloy formed by using the baseline Rene 80 composition andprocessed by direct metal laser melting (DMLM) process;

FIG. 3 illustrates a prior art microstructure image of an as-cast alloyformed by using the baseline Rene 80 composition;

FIG. 4 illustrates a high magnification microstructure image of anas-built alloy formed by using the baseline Rene 80 composition andprocessed by the DMLM process;

FIG. 5 illustrates a high magnification microstructure image of anas-built alloy formed by using the baseline Rene 80 composition andprocessed by the DMLM process;

FIG. 6 illustrates a low magnification microstructure image of a nickelbase super alloy formed by using the baseline Rene 80 composition andprocessed by the DMLM process; and

FIG. 7 illustrates a low magnification microstructure of a Ni-base superalloy, processed by the DMLM process, according to an embodiment of thepresent technique.

DETAILED DESCRIPTION

The present invention is generally applicable to systems that includenickel base (Ni-base) super alloys and methods of forming Ni-base superalloys that operate within environments characterized by relatively hightemperatures, and are therefore subjected to a hostile oxidizingenvironment and severe mechanical stresses. Notable examples of suchcomponents include turbine nozzles and blades, shrouds and augmentorhardware of gas turbine engines. While the advantages of this inventionwill be described with reference to gas turbine engine hardware, theteachings of the invention are generally applicable to any componentthat can be used in high temperature and harsh environments.

Approximating language, as used herein throughout the specification andclaims, may be applied to modify any quantitative representation thatcould permissibly vary without resulting in a change in the basicfunction to which it is related. Accordingly, a value modified by a termor terms, such as “about”, and “substantially” is not to be limited tothe precise value specified. In some instances, the approximatinglanguage may correspond to the precision of an instrument for measuringthe value. Here and throughout the specification and claims, rangelimitations may be combined and/or interchanged, such ranges areidentified and include all the sub-ranges contained therein unlesscontext or language indicates otherwise.

In the following specification and the claims, the singular forms “a”,“an” and “the” include plural referents unless the context clearlydictates otherwise. As used herein, the term “or” is not meant to beexclusive and refers to at least one of the referenced components beingpresent and includes instances in which a combination of the referencedcomponents may be present, unless the context clearly dictatesotherwise.

As used herein, the terms “may” and “may be” indicate a possibility ofan occurrence within a set of circumstances; a possession of a specifiedproperty, characteristic or function; and/or qualify another verb byexpressing one or more of an ability, capability, or possibilityassociated with the qualified verb. Accordingly, usage of “may” and “maybe” indicates that a modified term is apparently appropriate, capable,or suitable for an indicated capacity, function, or usage, while takinginto account that in some circumstances, the modified term may sometimesnot be appropriate, capable, or suitable.

Embodiments of the invention described herein address the notedshortcomings of the state of the art. Some embodiments present a methodof making intermediate alloys and Ni-base super alloys using additivemanufacturing techniques. Further, some other embodiments present amethod of making texture-free Ni-base super alloys using the additivemanufacturing techniques.

Additive manufacturing refers to a process by which digital threedimensional (3D) design data is used to build up a component byadding-in layers of material deposition. A material may be used in apowder form for building a component in a layer by layer manner.Additive manufacturing may include 3D printing, rapid prototyping (RP),direct digital manufacturing (DDM), layered manufacturing, and additivefabrication. Advantageously additive manufacturing is a design-drivenmanufacturing process that facilitates manufacturing of structureshaving complex design. Further, additive manufacturing provides a highdegree of design freedom, optimization and integration of functionalfeatures, and a relatively high degree of product customization.

Additive manufacturing may include certain specific processes, such asfor example, selective laser sintering, direct metal laser sintering,selective laser melting, and direct metal laser melting etc. While theembodiments disclosed herein are described with reference to directmetal laser melting (DMLM) process, other additive manufacturingtechniques along with required design and process variations may be usedfor producing an intermediate alloy, a Ni-base super alloy, or both theintermediate alloy and the Ni-base super alloy disclosed herein.

In certain embodiments, the DMLM process starts by applying a thin layerof a powder material to a building platform. A laser beam is used tomelt or fuse the powder at one or more defined portions. In one example,the portions may be defined by computer-generated component design data.Subsequently, a second layer of powder is applied on the previous layerof the powder. Optionally, the building platform may be adjusted (forexample, lowered) before applying the second layer of powder. Further,material in the second layer of powder may be melted or fused so as tobond the material in the second layer of powder with the layer below atone or more predefined portions. Similarly, subsequent layers of powdermay be deposited on the second layer and one or more portions in thesesubsequent layers may be melted and solidified to form bonds betweenadjacent layers. Further, the melted parts in intermediate steps orafter laying out all the layers results may be solidified in a desiredcomponent of required size and shape. Moreover, in some embodiments, theresulting components may be subjected to further heat-treatment toimpart desirable properties to the component, such as for example,required microstructure and high temperature stability.

A microstructure of a Ni-base super alloy may depend on the compositionof the super alloy. Consequently, during service, properties exhibitedby a component made from this Ni-base super alloy depend on thecomposition of the super alloy. Further, during service, propertiesexhibited by the component made from this Ni-base super alloy alsodepend on a method of making the super alloy during formation of thecomponent. Particularly, in the Ni-base super alloys, the method ofmanufacture and the further heat-treatment given to the component maydetermine the strength and robustness of the component.

In some embodiments, the DMLM process is used to build a component froma Ni-base super alloy. The Ni-base super alloy has a composition thatassists in achieving properties, such as, but not limited to, hightemperature strength, oxidation resistance, and creep resistance thatare desirable for high temperature and/or high pressure applications.One example of a Ni-base super alloy composition that may be used for agas turbine component is a Rene 80™ composition. A standard Rene 80composition may include cobalt (Co), chromium (Cr), molybdenum (Mo),tungsten (W), aluminum (Al), titanium (Ti), zirconium (Zr), carbon (C),and boron (B), along with nickel (Ni). Depending on requiredapplications, amounts of each of these elements may vary to a desirableextent in a given standard Rene 80 alloy. In some embodiments, astandard Rene 80 composition including about 5 to 15 wt. % of Co, 10 to20 wt. % of Cr, 3 to 6 wt. % of Mo, 3 to 6 wt. % of W, 2 to 4 wt. % ofAl, 4.8 to 5.2 wt. % of Ti, 0.01 to 0.05 wt. % of Zr, 0.15 to 0.2 wt. %of C, 0.001 to 0.030 wt. % of B is used for a given application for usein a gas turbine component. This standard Rene 80 composition may bereferred to as a “baseline Rene 80 composition” henceforth, in thisapplication.

Generally a Ni-base super alloy formed by using the baseline Rene 80composition includes a matrix comprising a gamma (γ) phase. Thegamma-phase is a solid solution with a face-centered cubic (fcc) latticeand randomly distributed different species of atoms. In someembodiments, the Ni-base super alloy further includes precipitates of agamma-prime (γ′) phase and/or gamma-double prime (γ″) phase.

In some embodiments, when a Ni-base super alloy having baseline Rene 80composition is processed using the DMLM process and subjected toheat-treatment post processing, the super alloy is observed to have adifferent creep behavior as compared to a super alloy of the samecomposition formed by other processes, such as a casting process.Without being bound by any theory, the present inventors assign thisdifference in the creep behavior to a microstructural difference thatoccurs in the as-formed super alloy of the baseline Rene 80 compositionas a result of the DMLM process. Possibly because of high cooling ratesthat are achieved in the DMLM process, a microstructure formed when theNi-base super alloy is made via the DMLM process is distinctly differentthan a microstructure that is formed when the alloy is made usingtraditional casting or investment casting processes.

FIG. 1 illustrates a low magnification prior art microstructure image ofan as-cast alloy 100 of baseline Rene 80 composition and FIG. 2illustrates a low magnification microstructure of an as-built alloy 200of baseline Rene 80 composition processed by DMLM method. As usedherein, the term “as-cast alloy” refers to an alloy made usingconventional methods of casting, and the term “as-built alloy” refers toan alloy made using the DMLM process. The microstructure of the as-builtalloy 200 of baseline Rene 80 composition includes a columnar structure.However, in the illustrated embodiment of FIG. 2, the columnar structureis too small to resolve into any finer details at the same magnificationas that illustrated in the as-cast alloy 100 of FIG. 1. FIG. 3 showsanother image of the prior art as-cast alloy 100 of baseline Rene 80composition with a magnification scale of 20 micrometers. FIGS. 4 and 5show comparatively high magnification images (than FIG. 2) of theas-built alloy 200 of the baseline Rene 80 composition. Specifically,FIGS. 4 and 5 illustrate the as-built alloy 200 of baseline Rene 80composition with magnification scales of 1 micrometer (μm) and 500nanometers (nms), respectively.

Comparing the as-cast alloy 100 of FIG. 3 with the as-built alloy 200 ofFIGS. 4 and 5, it can be seen that carbides 102 that are present in theas-cast alloy 100 of FIG. 3 are observed to be randomly scattered in thealloy 100, while the carbides 202 that are seen in the as-built alloy200 of FIGS. 4 and 5 are observed to be preferentially located in theintercolumnar regions 204. The intercolumnar regions 204 are defined bythe columnar regions 206 of the as-built alloy 200 that are present as asignature of the DMLM process.

Further, it can be noted that the carbides 102 that are present in themicrostructure of the as-cast alloy 100 of baseline Rene 80 compositionas seen in FIG. 3 are distinctly different in their morphology anddistribution than the carbides 202 that are present in themicrostructure of the as-built alloy 200 of baseline Rene 80 compositionas seen in FIGS. 4 and 5. For example, the carbides 102 of the as-castalloy 100 are relatively large in size as compared to the carbides 202that are present in the intercolumnar regions 204 of the as-built alloy200. For example, while the carbides 102 that are present in the as-castalloy 100 may have an average diameter of about 2-10 micrometers (μm),the average diameter of the carbides 202 that are observed in theas-built alloy 200 as seen in FIGS. 4 and 5 are less than about 300nanometers (nm).

Furthermore, it is observed by the inventors that the carbides 202 thatare present in the intercolumnar regions 204 of the as-built alloy ofbaseline Rene 80 composition are relatively more in number and formsubstantially closely spaced sections, such as films or arrays ofcarbides, in the intercolumnar regions 204. Formation of these films orarrays of smaller carbides in the intercolumnar regions 204 of theas-built alloy 200 was unexpected and were hitherto not noticed. By wayof example, in as-cast alloy 100, such substantially closely spacedsections are generally not found.

Generally, strengthening mechanisms of Ni-base super alloys are known tobe complex and mainly involve precipitation of intermetallic phases andcarbides in the grains as well as at the grain boundaries. There may bethree types of carbides which are found in Ni-base super alloys, namelyMC type, M₂₃C₆ type, and M₆C type carbides. MC type carbides arecarbides with MC composition, where M is a metal and C is carbon. MCtype carbides are known as primary carbides or solidification-typecarbides, and act as a source of carbon for secondary carbides (e.g.carbides of M₂₃C₆ and M₆C type). The primary and secondary carbides thatare present at the grain boundaries may hinder any movement ofdislocations and grain boundaries during formation or service of acomponent made by the Ni-base super alloys.

The high number density and close packing of the carbides 202 observedin intercolumnar regions 204 of the as-built alloy 200 may hinder anygrain growth during further heat-treatment of the as-built alloy 200.Moreover, the closely spaced carbides 202 may effectively limit stressrelaxation of grains through heat-treatment during formation or serviceof a component made by the as-built alloy 200. Additionally, failure toaccommodate stress build-up may result in crack formation in the alloy,thereby compromising mechanical integrity and high temperatureproperties of the component made by these alloys.

FIG. 6 illustrates a low magnification microstructure of a Ni-base superalloy 600 that is obtained after heat-treating the as-built alloy 200 ofthe baseline Rene 80 composition. The heat-treatment is carried out atabout 1200° C. using hot isotactic pressing (HIP) process. Comparing themicrostructure of the as-built alloy 200 of FIG. 2 with themicrostructure of the Ni-base super alloy 600 of FIG. 6, it is observedthat increase in grain sizes of the Ni-base super alloy 600 is notsignificant from the grain sizes of the as-built alloy 200. This lack ofgrain growth may be attributed to the well-pinned microstructure of theas-built alloy 200. In addition, the columnar nature of the grains, andsubsequent texture of the as-built alloy 200 is well-maintained in theNi-base super alloy 600 formed by this method. Accordingly, in certainembodiments, a Ni-base super alloy 600 of the baseline Rene 80composition, formed after heat-treatment of the as-built alloy 200obtained by DMLM process, may have substantially anisotropic mechanicalproperties, such as, for example, anisotropic creep behavior. In someembodiments, reduction in creep properties and directional variations increep capabilities of the Ni-base super alloy 600 may not be desirable.Without being bound by any particular theory, the inventors attributethis altered creep behavior to the reduced grain growth duringheat-treatment of the as-built alloy 200.

For some applications, such as, for example, in hot gas path components,a Ni-base super alloy microstructure that includes substantiallyequi-axed grains, a matrix of gamma-nickel solid-solution withgamma-prime and a fine dispersion of carbides is desired with thegrain-boundaries that are substantially free of embrittling carbidefilms or phases.

In some of these embodiments, the creep properties of the Ni-base superalloys 600 may be improved by enabling grain growth of the as-builtalloy 200 during heat-treatment. Further, in some embodiments, a graingrowth that yields a random orientation is desirable to more closelymatch the properties of the alloys built by the DMLM process with thatof the heat-treated, as-cast alloy.

Without being bound by any particular theory, the inventors envisagethat recrystallization and grain growth with substantially equi-axialorientation of the grains of the as-built alloy 200 may be formed byreducing the formation of the films or arrays of fine carbides in theintercolumnar regions of the alloys built by the DMLM process. Someembodiments described herein are directed towards reducing the carbidecontent in the intercolumnar regions of the as-built alloy.

In certain embodiments, carbide content in the grain boundaries may bereduced by using approaches, such as, but not limited to, reducing thecarbon content in an initial reactive powder, reducing a metal contentthat participates in carbide formation, enabling more primary carbideformation using the available carbon and hindering the secondary carbideformation, directing the available carbon to be reacted with otherelements than the material that are prone to form the secondary carbideformation, directing the carbon content to be in the grain regions, orcombinations thereof, thereby hindering the grain boundary segregationof carbides.

Some embodiments described herein are directed at the method of reducingthe carbide formation in the grain boundaries by reducing the carbonamount in the initial starting powder along with optional reduction inthe percentages of metallic elements that are involved in secondarycarbide formation. In certain embodiments, an amount of carbon presentin the initial starting powder considered for processing by DMLM processis reduced compared to generally used carbon content in a powder usedfor forming the Ni-base super alloy. For example, in a Ni-base superalloy having the baseline Rene 80 composition, if the carbon content isin a range from about 0.15 wt. % to about 0.2 wt. %, an alteredcomposition that is used in certain embodiments has a carbon contentthat is less than 0.15 wt. % of the overall powder composition.

In some embodiments, an amount of a primary carbide-forming element isalso reduced, along with reducing the carbon content in the startingpowder that is used for the formation of a Ni-base super alloy by theDMLM process. In some embodiments, an amount of titanium in the startingpowder is restricted along with restricting carbon element, to limit thecarbide formation in the alloy formed. For example, if the titanium isgenerally present in a range from about 4.8 wt. % to about 5.2 wt. % ina baseline Rene 80 powder composition used for casting, an altered Rene80 powder composition that is used in certain embodiments of the presentinvention has a titanium content that is less than 4.7 wt. % of theoverall powder composition.

Reducing an amount of a primary carbide-forming element, along withreducing carbon content, is particularly beneficial to retain a matrixalloy chemistry of the as-built alloy nominally similar to that of theas-cast alloy of the baseline Rene 80 composition. As used herein, the“matrix alloy chemistry” represents the alloy composition of the matrixphase. Retaining the matrix alloy chemistry is particularly advantageousto substantially avoid changes in some of the alloy chemistry-relatedproperties, such as, for example, phase instabilities, decrease ofstrength of the formed alloy, and lattice misfit between various phases.

In addition to or in place of reducing the carbon content in thestarting powder, in some embodiments, an amount of two or more primarycarbide-forming metallic elements are also reduced. In certainembodiments, total amount of primary carbide-forming metallic elementsis reduced to less than 5 wt. % of the overall composition of thestarting powder. This reduction in the primary carbide-forming metallicelements may be carried out along with or in place of reducing thecarbon content. As used herein, “primary carbide-forming metallicelements” include metallic elements that may form primary carbides inthe normal processing conditions of forming an as-built alloy using theDMLM process, and where the carbides thus formed are likely to segregateto the intercolumnar regions of the as-built alloy. In some embodiments,the primary carbide-forming metallic elements may include transitionmetal elements of group 4A and group 5A of the periodic table.Non-limiting examples for the primary carbide-forming metallic elementsmay include one or more of titanium, zirconium, hafnium, vanadium,niobium, and tantalum. In some embodiments, reduction in an amount ofthe primary carbide-forming metallic elements is in addition to thereduction in the carbon level in the starting powder, and may be inproportion to the reduction in the carbon level. In one embodiment, thereduction in a total atomic percentage of the primary carbide-formingmetallic elements is substantially equal to the atomic percentagereduction of carbon in the starting powder.

Further, without being bound by any theory, it is believed by theinventors that when a percentage of carbon and carbide-forming metallicelements are below a determined value in the initial powders that areused for forming the Rene 80 alloy with altered Rene 80 composition, theprimary carbides (of MC type) that are formed in the alloy are primarilyprecipitated from the gamma phase matrix on cooling aftersolidification, rather than during solidification from the melt itself.This is in contrast to forming an alloy with the baseline Rene 80composition, where the carbon and carbide-forming metals are present ina relatively high amount and induce the formation of solidification-type(i.e. solidifying from the melt) primary carbide precipitation. Thesesolidification-type carbides formed in the baseline Rene 80 compositionare observed to be located in and around the grain boundaries. However,the primary carbides that precipitate from the gamma phase matrix in thealtered Rene 80 composition are not confined to grain boundaries.Advantageously, the primary carbides that precipitate from the gammaphase matrix are well-distributed throughout the microstructure of theas-built alloy, though preferentially distributed in the intercolumnarregions of dendrites. Consequently, when an as-built alloy formed fromthe DMLM process using altered Rene 80 composition (having loweredamounts of carbon and carbide-forming elements) is heat-treated, theprimary carbides that are well-dispersed in the microstructure mayre-precipitate into the secondary carbides (M₂₃C₆/M₆C) during subsequentheat treatment, and may be well-dispersed both in the grains and in thegrain boundaries of the heat treated Ni-base super alloy thus formed.

In some embodiments, a method for making the as-built alloy using theDMLM process includes at least partially melting and solidifying apowder that includes about 5 to 15 wt. % of Co, 10 to 20 wt. % of Cr, 3to 6 wt. % of Mo, 3 to 6 wt. % of W, 2 to 4 wt. % of Al, 4.2 to 4.7 wt.% of Ti, 0.01 to 0.05 wt. % of Zr, 0.015 to 0.060 wt. % of C, and 0.001to 0.030 wt. % of B and balance substantially Ni. In certainembodiments, the titanium content may be in a range from about 4.2 wt. %to about 4.7 wt. % of the powder, and in some further embodiments, thetitanium content may be in a range from about 4.4 wt. % to about 4.6 wt.% of the powder. Further, in certain embodiments, the carbon content maybe in a range from about 0.01 wt. % to about 0.04 wt. % of the powder,and in some further embodiments, the titanium content may be in a rangefrom about 0.01 wt. % to about 0.03 wt. % of the powder. The selectedcarbon ranges provided herein particularly aid in reducing densedistribution of carbides in the grain and intercellular boundaries.

In some embodiments, an as-built alloy is formed by using a powder withthe altered Rene 80 composition having carbon in a range from about0.015 wt. % to about 0.060 wt. % and titanium in an amount in a rangefrom about 4.2 wt. % to about 4.7 wt. % of the powder. This as-builtalloy, formed by using the powder having the above-mentioned alteredRene 80 composition may be referred to as an “intermediate alloy”henceforth, in this application.

In some embodiments, partial melting and solidification may be carriedout as part of the DMLM process in building the intermediate alloy. Asdescribed hereinabove, the intermediate alloy formed by melting andsolidifying the altered Rene 80 composition has columnar regions. Acolumnar region is a region of elongated dendrites having a preferredorientation. Columnar regions may be formed as a result of competitivegrowth at a particular direction during the solidification step.Intercolumnar regions are the regions between adjacent columnar regions.Dendritic growth is a general form of crystal growth encountered whenmetals, alloys and many other materials solidify under low thermalgradients. The dendritic growth includes columnar regions andintercolumnar regions.

A dendrite or dendrite structure is generally characterized by themicrostructure parameters associated with the dendrites. Themicrostructure of the dendrite structure of the intermediate alloy isnormally characterized by primary dendrite or cellular arm spacing. Aprimary dendrite arm spacing is a characteristic length scale that isused to determine segregation patter in in alloy solidification.Generally the primary dendrite arm spacing is obtained by measuringdistances between cores (centers) of neighboring dendrites. In someembodiments, a primary dendrite arm spacing of the dendrite structure ofthe intermediate alloy is less than about 3 micrometers. In certainembodiments, the primary dendrite arm spacing of the dendrite structureof the intermediate alloy is less than about 2 micrometers. In certainembodiments, a primary dendrite arm spacing of the dendrite structure isless than about 3 micrometers. The intermediate alloy has a compositionthat includes about 5 to 15 wt. % of Co, 10 to 20 wt. % of Cr, 3 to 6wt. % of Mo, 3 to 6 wt. % of W, 2 to 4 wt. % of Al, 4.2 to 4.7 wt. % ofTi, 0.01 to 0.05 wt. % of Zr, 0.015 to 0.060 wt. % of C, 0.001 to 0.030wt. % of B and balance substantially Ni.

In some embodiments, the dendrite structure that is formed in theintermediate alloy has only primary arms and arm spacings and may notcontain any substantial secondary arms and secondary arm spacings. Thesedendrite structures that only have the primary arms and arm spacings arereferred to as cell structures.

During melting and solidification of the powder, multiple chemicalreactions may occur between the elements present in the powder to formthe intermediate alloy. Upon melting and solidification, theintermediate alloy, having relatively low levels of carbon andcarbide-forming metallic elements, may have relatively low level ofmetal carbide formation as compared to an as-cast alloy using thebaseline Rene 80 composition with regular levels of carbon andcarbide-forming metallic elements. Accordingly, in some embodiments ofthe method, an amount of the metal carbides present in the intermediatealloy is less than about 0.5 mole % of the composition of theintermediate alloy. In particular, the initial powder composition thatis used for the formation of the intermediate alloy includes carbideformers in an amount such that an equilibrium carbide content of theintermediate alloy is less than about 0.5 mole % at temperatures aboutthe solidus temperature. In certain embodiments, the amount of metalcarbides may be further limited to be less than about 0.3 mole % of theintermediate alloy composition.

In certain embodiments, the metal carbides that are precipitated in theintermediate alloy may be present both in the columnar and intercolumnarregions. In some embodiments, the metal carbides are disposed in theintercolumnar regions of the intermediate alloy. Metal carbides presentin the intercolumnar regions may hinder the grain growth to a greaterextent during a heat-treatment as compared to metal carbides that arepresent within the columnar regions. In some embodiments, an amount ofthe metal carbides present in the intercolumnar regions of the dendritestructure is less than about 0.3 mole % of the intermediate alloycomposition. The amount of metal carbides in the intercolumnar regionsmay further be limited to be less than about 0.2 mole % of theintermediate alloy composition, in accordance to certain embodiments.

The method further includes heat-treating the intermediate alloy in atemperature range from about 1050° C. to about 1250° C. to form atexture-free Ni-base super alloy. This Ni-base super alloy, formed byheat-treating the intermediate alloy (of altered Rene 80 composition)may be referred to as a “modified Ni-base super alloy” henceforth, inthis application.

It may be noted that composition of the modified Ni-base super alloy isconsidered to have substantially similar composition as that of theintermediate alloy. In particular, there is no substantial change thatoccurs in the matrix alloy chemistry of the intermediate alloy, when theintermediate alloy is subjected to the heat-treatment to form themodified Ni-base super alloy. The matrix alloy composition of themodified Ni-base super alloy is substantially formed in the intermediatealloy form itself. Specifically, the heat-treatment step used forconverting the intermediate alloy to the modified Ni-base super alloy isa processing step that is used for altering the microstructure of theintermediate alloy to that of the modified Ni-base super alloy, than forforming the composition of the modified Ni-base super alloy. A slightvariation in the composition of the intermediate alloy and the modifiedNi-base super alloy, if present, may be due to an interaction betweenthe already formed intermediate alloy with the environment. Any suchvariation in the composition during heat-treatment is limited to lessthan about 1 volume % of the modified Ni-base super alloy. In someembodiments, the method includes heat-treating the intermediate alloy ina temperature range from about 1150° C. to about 1250° C. to form themodified Ni-base super alloy.

The heat-treatment that is imparted to an intermediate alloy having alow carbide content in the intercolumnar regions as noted above allowsfor a substantial grain growth in the modified Ni-base super alloy thusformed. The modified Ni-base super alloy includes grains, grainboundaries, and metal carbides. The metal carbides may be disposed inthe grains or the grain boundaries. In some embodiments, the modifiedNi-base super alloy is substantially free of carbides that areprecipitated from the melt or liquid state of the alloy. In someembodiments, the modified Ni-base super alloy includes metal carbidesthat are precipitated from the solid gamma phase matrix. Further, anequi-axial orientation is observed in the Ni-base super alloy obtainedfrom the powder with modified (lowered) carbon content. Hence thelowered carbon content aids to achieve a substantially modified Ni-basesuper alloy.

FIG. 7 illustrates a low magnification microstructure of a modifiedNi-base super alloy 700 that is obtained after heat-treating anintermediate alloy. The heat-treatment is carried out at about 1200° C.In some embodiments, heat-treatment is carried out using hot isotacticpressing (HIP) process. While the heat-treatment disclosed above isparticularly directed to form the desired microstructural change in theintermediate alloy to form the modified Ni-base super-alloy, there maybe certain other involuntary changes that may occur in the alloy duringheat-treatment, such as, for example, change in the gamma-prime phasedistribution. In some embodiments, an additional heat-treatment may beimparted to the heat-treated intermediate alloy to form the modifiedNi-base super alloy 700. By way of example, the intermediate alloy maybe subjected to the additional treatment to obtain carbidere-precipitation and/or to achieve a preferred gamma-prime distribution.

An effect of the amount of carbon or carbide formers in the initialpowder can be clearly observed by comparing the microstructure of themodified Ni-base super alloy 700 of FIG. 7 with that of a Ni-base superalloy 600 that is obtained after heat-treating an as-built alloy of thebaseline Rene 80 composition, as depicted in FIG. 6. The microstructureof the modified Ni-base super alloy 700 depicts larger and equi-axialgrains as compared to the columnar grains with lower grain sizes of theNi-base super alloy 600. The size and orientation of individual grainsin the modified Ni-base super alloy 700 were further measured throughthe use of electron backscatter detection (EBSD) techniques (not shown).EBSD showed that the modified Ni-base super alloy 700 resulted inrandomly oriented grains after heat treatment as well as substantialgrain growth over the Ni-base super alloy 600 obtained by the baselineRene 80 composition.

Thus in some embodiments, the modified Ni-base super alloy 700 formed bythe above-described method is texture-free. As used herein the“texture-free alloy” is defined as “the alloy, wherein a preferredcrystallographic orientation in any direction is less than 20% by volumeof any representative region of the alloy considered”. In someembodiments, the crystallographic orientation in the modified Ni-basesuper alloy 700 is less than 10 volume %, and in certain embodiments,the modified Ni-base super alloy 700 is substantially free of apreferred crystallographic orientation in any particular direction.

Considering that the composition of the modified Ni-base super alloy issubstantially the same as the composition of the intermediate alloy, insome embodiments, an amount of the metal carbides present in thecomposition of the modified Ni-base super alloy is less than about 0.5mole % of the composition. The amount of metal carbides may further belimited to be less than about 0.3 mole % of the modified Ni-base superalloy composition, in accordance to certain embodiments. The amount ofmetal carbides in the modified Ni-base super alloy may further belimited to be less than about 0.2 mole % of the modified Ni-base superalloy composition, in accordance to certain embodiments. Furthermore, anaverage size of the metal carbides in the modified Ni-base super alloymay be less than about 1 micrometer.

In some embodiments, the modified Ni-base super alloy formed hereinincludes a composition that includes about 5 to 15 wt. % of Co, 10 to 20wt. % of Cr, 3 to 6 wt. % of Mo, 3 to 6 wt. % of W, 2 to 4 wt. % of Al,4.2 to 4.7 wt. % of Ti, 0.01 to 0.05 wt. % of Zr, 0.015 to 0.060 wt. %of C, 0.001 to 0.030 wt. % of B and balance substantially Ni. Themodified Ni-base super alloy may be formed by the method describedhereinabove and has a texture-free form that includes a gamma phasematrix, a precipitated gamma-prime phase, and metal carbidesprecipitated from gamma phase matrix. The metal carbides that aredisposed in the modified Ni-base super alloy is less than about 0.3 mole% of the composition and an average size of the metal carbides that arepresent in the modified Ni-base super alloy is less than about 1micrometer.

Advantageously, the modified Ni-base super alloy has a reduced carbidecontent in its composition, than just a reduced carbon amount. Thereduced carbide content aids to maintain a local chemistry of thedifferent phases of (other than carbides) present in the alloy, such as,for example, gamma and gamma-prime phases. Maintaining the localchemistry is particularly advantageous in that the alloy is notsusceptible to size alterations and changes in lattice misfit betweenthe matrix and precipitating phases.

Further, reduction of carbide content (rather than carbon content) inthe modified Ni-base super alloy and maintaining the local phasechemistry similar to a standard parent alloy prepared using traditionalroutes, enhances long term microstructure stability of the modifiedNi-base super alloy and further supports joining the modified Ni-basesuper alloy formed to an existing standard alloy.

While only certain features of the invention have been illustrated anddescribed herein, many modifications and changes will occur to thoseskilled in the art. It is, therefore, to be understood that the appendedclaims are intended to cover all such modifications and changes as fallwithin the true spirit of the invention.

The invention claimed is:
 1. An alloy, comprising: a dendrite structurecomprising columnar regions and intercolumnar regions, wherein a primarydendrite arm spacing of the dendrite structure is less than about 3micrometers; and a composition comprising about 5 to 15 wt % of Co, 10to 20 wt % of Cr, 3 to 6 wt % of Mo, 3 to 6 wt % of W, 2 to 4 wt % ofAl, 4.2 to 4.7 wt % of Ti, 0.01 to 0.05 wt % of Zr, 0.015 to 0.060 wt %of C, 0.001 to 0.030 wt % of B and balance substantially Ni.
 2. Thealloy of claim 1, wherein an amount of metal carbides present in thealloy is less than about 0.5 mole % of the composition.
 3. The alloy ofclaim 2, wherein an amount of the metal carbides present in the alloy isless than about 0.3 mole % of the composition.
 4. The alloy of claim 2,wherein the metal carbides are primarily disposed in the intercolumnarregions of the dendrite structure.
 5. The alloy of claim 2, wherein anaverage diameter of the metal carbides is less than 300 nanometers. 6.The alloy of claim 1 formed by a direct metal laser melting (DMLM)method.
 7. A texture-free Ni-base super alloy, comprising: a compositioncomprising 5 to 15 wt % of Co, 10 to 20 wt % of Cr, 3 to 6 wt % of Mo, 3to 6 wt % of W, 2 to 4 wt % of Al, 4.2 to 4.7 wt % of Ti, 0.01 to 0.05wt % of Zr, 0.015 to 0.060 wt % of C, 0.001 to 0.030 wt % of B andbalance substantially Ni; a gamma phase matrix; precipitates ofgamma-prime phase; and metal carbides precipitated from the gamma phasematrix, wherein an amount of the metal carbides present in the Ni-basesuper alloy is less than about 0.3 mole % of the Ni-base super alloy;and an average size of the metal carbides is less than about 1micrometer.
 8. The texture-free Ni-base super alloy of claim 7, whereinan amount of the gamma-prime phase in the Ni-base super alloy is greaterthan 40 volume percent.
 9. The texture-free Ni-base super alloy of claim7 formed by a direct metal laser melting (DMLM) method and heat-treatingat a temperature range from about 1050° C. to about 1250° C.
 10. Thetexture-free Ni-base super alloy of claim 9, wherein the Ni-base superalloy is substantially free of metal carbides precipitated from meltduring solidification step of the DMLM method.